1. Field of the Invention
The present invention relates generally to compositions of matter and methods for synthesizing a composition of matter including controlling the pore size, pore size distribution and porosity of aluminum-oxide based ceramics through the choice of substituents on carboxylate-alumoxanes and aluminum-oxide nanoparticles. The invention includes aluminum and aluminum oxide ceramic bodies with intra-granular pores in the nanometer range and methods for forming intra-granular pores in the nanometer range in alumina and aluminum oxide ceramic bodies. The invention provides for the control over pore size and pore size distribution by the use of chemical substituents on the carboxylate-alumoxanes and aluminum-oxide nanoparticles. The invention also includes the use of controlled-porosity ceramics for ceramic membrane filters and coatings and interphase layers for fibers and fiber reinforced composites.
2. Description of the Related Art
The oxides and hydroxides of aluminum are undoubtedly among the most industrially important chemicals. Their uses include: precursors for the production of aluminum metal, catalysts and absorbents; structural ceramic materials; reinforcing agents for plastics and rubbers, antacids and binders for the pharmaceutical industry; and as low dielectric loss insulators in the electronics industry. Traditional ceramic processing involves three basic steps generally referred to as powder-processing, shape-forming, and densification, often with a final mechanical finishing step (Kingery et al. 1976 and Richerson 1992). Whereas traditional sintering process are primarily for the manufacture of dense parts, solution-gelation processes have been applied industrially used for the production of porous materials and coatings. Solution-gelation involves a four stage process: dispersion; gelation; drying; and firing. A stable liquid dispersion or sol of the colloidal ceramic precursor is initially formed in a solvent with appropriate additives. By change in the concentration (aging) or the pH, the dispersion is polymerized to form a solid dispersion or gel. The excess liquid is removed from this gel by drying, and the final ceramic is formed by firing the gel at higher temperatures. The common solution-gelation route to aluminum oxides employs aluminum hydroxide (or hydroxide-based material) as the solid colloid, with the second phase being water and/or an organic solvent. Aluminum hydroxide gels have traditionally been prepared by the neutralization of a concentrated aluminum salt solution (Serna et al. 1977), however, the strong interactions of the freshly precipitated alumina gels with ions from the precursor solutions makes it difficult to prepare these gels in pure form (Green and Hem 1974). To avoid this complication alumina gels may be prepared from the hydrolysis of aluminum alkoxides, Al(OR)3 (Eq. 1). 
Although this method was originally reported by Adkins in 1922, it was not until Teichmer et al. (1976) reported the preparation of alumina aerogels, and Yoldas (1975) showed that transparent ceramic bodies can be obtained by the pyrolysis of suitable alumina gels, that interest increased significantly. Other pertinent references include: Nogami (1994), Low et al. (1997), Nikolic and Radonjic (1997), Rezgui and Gates (1997), Rezgui et al. (1994). The exact composition of the gel in commercial systems is ordinarily proprietary, however, a typical composition will include an aluminum compound, a mineral acid and a complexing agent to inhibit premature precipitation of the gel. The aluminum compound has traditionally been the direct precursor to pseudo-boehmite.
The aluminum based sol-gels formed during the hydrolysis of aluminum compounds belong to a general class of compounds, namely alumoxanes. These materials were first reported in 1958 (Andrianov and Zhadanov, 1958) with siloxide substituents, however, they have since been prepared with a wide variety of substituents on aluminum. Recent work has shown that the structure of alumoxanes is as three dimensional cage compounds (Apblett et al. 1992 and Landry et al. 1993). For example, siloxy-alumoxanes, [Al(O)(OH)x(OSiR3)1−x]n, consist of an aluminum-oxygen core structure (FIG. 1) analogous to that found in the mineral boehmite, [Al(O)(OH)]n, with a siloxide substituents. In the siloxy-alumoxanes, the“organic” is typically like that shown in FIG. 2. However, the carboxylate anion, [RCO2]−, is an isoelectronic and structural analog of the organic portion found in the siloxy-alumoxanes (FIG. 3). Based upon this approach the reaction of boehmite, [Al(O)(OH)]n, with carboxylic acids, has been developed (Landry et al. 1995) or Eq. 2. Carboxylate-substituted alumoxanes have been well characterized (Landry et al. 1995 and Callender et al. 1997). Solution particle-size measurements shows that carboxylate-alumoxanes are nano-particles with sizes ordinarily ranging from 1-1000 nm (FIG. 10, 11 and 12). Nano-particles are ordinarily defined as materials with sizes ranging from 1 nm to 1 μm. The carboxylate ligand is bound to the aluminum surface, and is only removed under extreme conditions. The carboxylate-alumoxane materials prepared from the reaction of boehmite and carboxylic acids are air and water stable materials and are easily processable (FIG. 7). The soluble carboxylate-alumoxanes can be dip-coated, spin coated, and spray-coated onto various substrates. The physical properties of these alumoxanes are highly dependent on the identity of the alkyl substituents, R, and range from those associated with insoluble crystalline powders to powders that readily form solutions or gels in hydrocarbon solvents and/or water. These alumoxanes are indefinitely stable under ambient conditions, and are adaptable to a wide range of processing techniques. The alumoxanes can be easily converted to aluminum oxide upon mild thermolysis, while they also react with metal complexes to form doped or mixed aluminum oxides (Kareiva et al. 1996).
The control of porosity (pore size, pore size distribution and pore density) is an important aspect of ceramics. Lower porosity improves strength, load-bearing capacity, and corrosion resistance, but can also lead to catastrophic failure from thermal shock, because the pores present act as crack stoppers in more porous ceramics. Cracks propagate intergranularly (between grains) and therefore the grain boundary toughness plays a large role in determining the fracture mode. Porosity between grains can promote crack propagation and lower the strength of a ceramic body. In both traditional and sol-gel processes, the porosity of the resulting ceramic is controlled through physical processing variables (Wilson and Stacey, 1981), such as the time or temperature of firing and the addition of pre-fired additives to seed crystal growth (Shelleman et al. 1986). Direct chemical control has not been observed. Furthermore, the pore size, pore size distribution and porosity are functions of the ceramic particles used to make the ceramic body, because the porosity is determined by the gaps between the individual particles (FIG. 4) and is therefore inter-granular, that is between the crystal grains. For example, pores below 0.1 μm in diameter require that submicron powders be used (in traditional ceramic processing), while smaller pores require sol-gel processing.
A particularly important area where the strength and porosity of ceramic materials is affected by the formation of inter-phase materials in fiber reinforced ceramic matrix and metal matrix composites. Fiber reinforced ceramic matrix composites (FRCMCs) are potential candidates for use in high temperature structural applications (Courtright, 1991). For example, aerospace applications include high thrust-to-weight ratio gas turbine engines and high-specific-impulse rocket motors. Ground based applications include, high efficiency turbine and diesel engines. In each of these applications there is a need for high performance ceramic materials that can be readily fabricated into complex shapes. Compared to current materials (e.g., nickel based superalloys) and proposed metallic and intermetallic matrix composites, FRCMCs have higher strengths at lower densities, higher maximum use temperatures, and better oxidation resistance. Ceramic materials are well known for their stability at high temperatures, adequate strength and resistance to corrosion, and can meet most of the requirements for gas turbine applications. However, the brittle nature of ceramic materials and their tendency to undergo catastrophic failure has limited their usefulness. By reinforcing ceramic materials with fibers, catastrophic failures can be reduced or eliminated. A major drawback in existing fiber reinforced ceramic matrix composites (FRCMCs) is the absence of a fiber-matrix interface (or interphase) that is weak and stable over the entire range of expected use. Limitations of such prior art FRCMCs are the instability of known interfaces and the chemical reactivity of many weak interphases with the fiber and/or matrix. The chemical design of interfaces to optimize the adhesion or transfer of load between reinforcing phase and the matrix, to enhance crack deflection through debonding or to control interfacial reactivity/stability are an important development. For both the fiber and ceramic matrix, material requirements include: high melting points, high modulus, low density, freedom from destructive phase transformations, low volatility, oxidative stability, and creep resistance. For structural applications at high temperatures, environmental stability and creep resistance are the dominant factors in determining the usefulness of ceramic materials. In general, monolithic polycrystalline oxide ceramics lose strength above 1200° C. Therefore, monolithic ceramics must be strengthened with high modulus fibers. The only materials that retain strength at these high temperatures, and under severe oxidative environments, are oxide fibers (e.g., sapphire) or silicon carbide (SiC) fiber. An additional concern is that the matrix and fiber materials must be chemically compatible (i.e., not react with each other). In fiber reinforced ceramic, the reinforcement is to enhance the fracture toughness. The fiber reinforcement prevents catastrophic brittle failure by providing mechanisms to dissipate energy during fracture. The operation of various toughening mechanisms, such as crack deflection, fiber pull out, and fiber bridging, depend to a large extent on the degree of chemical and/or mechanical bonding at the fiber-matrix interface. This chemical bonding is affected by the fiber surface chemistry and chemical reactivity between the fiber and matrix. The mechanical bonding is primarily controlled by the fiber surface morphology and the fiber/matrix thermal expansion match. In general, composites with strong interfacial bonding exhibit brittle behavior, characterized by high strength and low fracture toughness. If the interfacial interaction is weak then a composite will fail by catastrophic manner, and show high fracture toughness but low strength. It is therefore highly desirable to control the interfacial bond in order to optimize the overall mechanical behavior of the composite. The fiber-matrix interface must be sufficiently weak to allow debonding and sliding when a crack impinges upon it from the matrix; otherwise the crack passes through the fiber (or the fiber fails near the crack tip) and there is minimal or no toughening (Michalke and Hellmann, 1988). To control the strength of fiber coatings and the interaction between the coating and both the fiber and matrix, is extremely important to control the porosity of the coating materials.
In contrast, control of pore size, pore size distribution and porosity in ceramics is important for their applications in ceramic membranes and catalyst supports. Membrane-based technologies play a unique and increasingly important role in pollution prevention, resource recovery and waste treatment activities (Baker, 1991). Due in large part to cost considerations, polymeric membranes have dominated these environmental separations applications. However, the use of polymeric membranes in separations involving aggressive materials such as many solvents, acids, bases, and oxidants may be limited by the tolerance of these membranes to extreme conditions (Hsieh, 1988). Ceramic membranes are noted for their excellent mechanical strength and tolerance to solvents, pH, oxidant, and temperature extremes. In addition, the amphoteric properties of ceramic membrane surfaces result in a uniquely versatile membranes for water and waste water treatment. Membrane charge, selectivity, and permeation rate vary as a function of pH, ionic strength and other characteristics of the feed water solution chemistry, Baltus (1997) and Kim and Stevens (1997). Membrane characteristics as well as the properties of the contaminants can be manipulated through adjustments in the solution chemistry of the feed stream in one or more pretreatment steps (Anderson et al. 1988). Ceramic membranes are typically produced by slip casting a colloidal suspension on a porous ceramic support: Okubo, et al. (1990), Elaloui et al. (1997), Lin et al. (1991), Lao et al. (1994), Zaspalis et al. (1992). A schematic view of a typical membrane design is shown in FIG. 5. The individual membranes are mounted into a membrane module (see FIG. 6). Control of the colloidal suspension in the sol-gel process and limitations on the size of colloids that can be produced have constrained the range of membrane types that can be produced. In addition, a key obstacle to overcome in advancing the use of ceramic membranes for pollution prevention applications is cost. The sol-gel processes currently used to produce commercially available ceramic membranes is energy intensive and difficult to control. Considerable time and expense is invested in verifying membrane integrity and re-casting membranes to repair imperfections. Alternative approaches for manufacture of ceramic membranes include the anodic oxidization of aluminum metal membranes (Furneaux et al. 1989), pore size being controlled by the applied voltage used in the anodic oxidation. However, strong dielectric solutions of various acids must be employed, and ion beam or chemical etching is performed to produce a working filter. An ideal ceramic membrane must be highly selective, highly permeable, and highly durable. The membrane selectivity is primarily dependent upon the pore-size distribution: a narrow distribution contributes to a highly selective membrane. Membrane permeability is a function of global porosity, membrane thickness, connectivity, and pore-size distribution. Membrane durability is obtained by high homogeneity and high density; the latter entails a clear compromise with permeability. Mechanical integrity is enhanced in such application by slip-casting a relatively thin selective membrane onto a larger, durable membrane of poor selectivity but high permeability.